High efficiency visible and ultraviolet nanowire emitters

ABSTRACT

GaN-based nanowire heterostructures have been intensively studied for applications in light emitting diodes (LEDs), lasers, solar cells and solar fuel devices. Surface charge properties play a dominant role on the device performance and have been addressed within the prior art by use of a relatively thick large bandgap AlGaN shell covering the surfaces of axial InGaN nanowire LED heterostructures has been explored and shown substantial promise in reducing surface recombination leading to improved carrier injection efficiency and output power. However, these lead to increased complexity in device design, growth and fabrication processes thereby reducing yield/performance and increasing costs for devices. Accordingly, there are taught self-organising InGaN/AlGaN core-shell quaternary nanowire heterostructures wherein the In-rich core and Al-rich shell spontaneously form during the growth process.

CROSS-REFERENCE TO RELATED APPLICATIONS

This application is a divisional (continuation) application of U.S.patent application Ser. No. 15/177,608, filed Jun. 9, 2016, herebyincorporated by reference in its entirety, and which claims the benefitof priority from U.S. Provisional Patent Application Ser. No. 62/172,874filed Jun. 9, 2015, entitled “High Efficiency Visible and UltravioletNanowire Emitters.”

STATEMENT OF GOVERNMENT INTEREST

This invention was made with U.S. Government support under grantW911NF-12-1-0477 awarded by U.S. Army Research Office. The U.S.Government has certain rights in this invention.

FIELD OF THE INVENTION

This invention relates to semiconductor nanowire devices and moreparticularly to ultraviolet emitters, self-organized nanoclusters,Anderson localization, and self-organized elementally rich shells.

BACKGROUND OF THE INVENTION

Compound semiconductor nanowire structures allow broad spectrum photonicdevices to be implemented with a single growth process thereby greatlyincreasing the efficiency of devices exploiting them such as solar cellsand light emitting diodes (LEDs), for example, and reducing their costs.To date solid state emitters exploiting compound semiconductor nanowireshave been demonstrated in a range of materials for different wavelengthranges as depicted in Table 1. However, more recently improved growthtechniques, improved understanding of the effects affecting performance,the introduction of graded growth and self-organizing quantum structureshave allowed multiple emitters to be integrated within the same nanowireto provide white LEDs as well as increasing emission efficiency. Forexample, InGaN nanowires have been demonstrated from the red-amber sideof the spectrum through to the blue-violet side and near ultraviolet(UV) down to approximately 325-350 nm.

TABLE 1 Examples of Semiconductor Alloys for LEDs WavelengthSemiconductor System Mid-Infrared λ ≥1100 nm  InGaAsP Infrared λ ≥760 nmGaAs, AlGaAs Red 610 nm ≤ λ ≤ 760 nm AlGaAs, GaAsP, AlGaInP, GaP Orange590 nm ≤ λ ≤ 610 nm GaAsP, AlGaInP, GaP, InGaN, Yellow 570 nm ≤ λ ≤ 590nm GaAsP, AlGaInP, GaP, InGaN Green 500 nm ≤ λ ≤ 590 nm GaP, AlGaInP,AlGaP, InGaN, Blue 450 nm ≤ λ ≤ 500 nm Zinc selenide (ZnSe), InGaN, GaNViolet 400 nm ≤ λ ≤ 450 nm InGaN, GaN Ultraviolet λ ≤400 nm Boronnitride (BN), AlN, AlGaN, AlGaInN

Despite the progress in electrically injected semiconductor lasers inthe visible, infrared, and terahertz wavelength ranges it has remaineddifficult to realize electrically injected semiconductor lasers orefficient light emitting diodes (LEDs) within the rich deep ultraviolet(UV) spectrum 10-21. Bridging this deep UV gap would allow thereplacement of conventional mercury lamps by efficient solid-state UVlight sources for a broad range of applications, including waterpurification, disinfection, bio-chemical detection, medical diagnostics,and materials processing, to name a few. In this context, AlGaN-basedmaterials, with a direct energy bandgap in the range of 3.4 eV to 6.1eV, have been intensively studied.

However, whilst optically pumped AlGaN quantum well lasers UV-B and UV-Cbands have been demonstrated these have had relatively high thresholdsas a result of the properties of conventional AlGaN materials includingthe large bandgap and large effective mass for both electrons and holes.Reducing this can be achieved by modifying the density of states (DOS)using quantum-confined nanostructures, such as quantum dots and quantumwires. With the use of such low-dimensional nanostructures,semiconductor lasers with significantly enhanced gain and differentialgain have demonstrated. Quantum dot-like nanoclusters can also beinduced by phase separation where, for example, the presence of In-richnanoclusters has been commonly observed in InGaN-based quantum welllasers; and the resulting carrier localization has been identified asone of the major factors contributing to the excellent performance ofGaN-based quantum well lasers operating in the near-UV, blue, andblue-green spectral ranges. However, the relatively small latticemismatch (a maximum of 3% between GaN and AlN), has to date prohibitedthe realization of electrically injected quantum dot lasers in the deepUV band.

Accordingly, it would be beneficial to establish the formation ofself-organized Ga(Al)N quantum dots in the deep UV spectral rangeallowing low-dimensional quantum-confined nanostructures, such asquantum dots and quantum wires, to be achieved allowing deep UVsemiconductor lasers with significantly enhanced gain and differentialgain to be implemented.

GaN-based nanowire heterostructures have been intensively studied forapplications in light emitting diodes (LEDs), lasers, solar cells andsolar fuel devices. Recent studies have shown that the surface chargeproperties play a dominant role on the device performance such that forthe commonly reported axial nanowire LED designs they exhibit very lowoutput power as a result of the large surface recombination andresulting poor carrier injection efficiency. Radial variations of In/Gadistribution have been observed in InGaN/GaN dot/disk/well-in-a-wireheterostructures. However, such radial variations were found to beinsufficient to suppress non-radiative surface recombination underelectrical injection. In this regard, the use of a large bandgap AlGaNshell covering the surfaces of axial InGaN nanowire LED heterostructureshas been explored and shown substantial promise in reducing surfacerecombination leading to improved carrier injection efficiency andoutput power. In these approaches, however, either relatively thickAlGaN layers were grown either on the top p-GaN region of the InGaN/GaNnanowires or incorporated within the device active regions, In each casethe intention being to form an AlGaN shell for surface passivation.However, each approach leads to increased complexity in the devicedesign, growth and fabrication processes thereby reducingyield/performance and increasing costs for devices. Moreover, afundamental understanding of the impact of the AlGaN shell structure onthe carrier dynamics and device performance has remained elusive.

Accordingly, it would be beneficial to provide designers ofsemiconductor nanowire emitting devices and their manufacturingoperations with a means of implementing InGaN/AlGaN core-shellquaternary nanowire heterostructures wherein the In-rich core andAl-rich shell spontaneously form during the growth process. It would befurther beneficial for these core-shell quaternary nanowireheterostructures to be tunable in emission wavelength across the visiblespectral range allowing discrete high efficiency coloured nanowire LEDs,multi-colour high efficiency nanowire LEDs, and white high efficiencynanowire LEDs to be formed through adjustments in the growth parameters.Further, the inventors beneficially establish a direct correlationbetween the output power, carrier lifetime, and shell thickness toprovide a robust, large bandgap shell structure methodology fordramatically enhancing the performance of axial nanowire LEDs for thesolid state lighting and display applications.

Other aspects and features of the present invention will become apparentto those ordinarily skilled in the art upon review of the followingdescription of specific embodiments of the invention in conjunction withthe accompanying figures.

SUMMARY OF THE INVENTION

It is an object of the present invention to mitigate limitations in theprior art relating to relates to semiconductor nanowire devices and moreparticularly to ultraviolet emitters, self-organized nanoclusters,Anderson localization, and self-organized elementally rich shells.

In accordance with an embodiment of the invention there is provided amethod of forming atomic scale compositional modulations within acompound semiconductor nanowire to form three dimensional quantumconfinement of charge carriers by increasing the concentration of apredetermined element within the composition of the compoundsemiconductor above a threshold.

In accordance with an embodiment of the invention there is provided adevice comprising:

-   a plurality of nanowires grown on a substrate having diameters over    a predetermined region of their length within a predetermined range    of diameters and having a predetermined fill factor over a    predetermined region of the substrate;-   upper and lower electrical contacts formed on the plurality of    nanowires of the predetermined region of the substrate; wherein-   the plurality of nanowires coherently emit optical radiation through    the Anderson localization of light.

In accordance with an embodiment of the invention there is provided amethod comprising:

-   growing a plurality of nanowires on a substrate having diameters    over a predetermined region of their length within a predetermined    range of diameters and having a predetermined fill factor over a    predetermined region of the substrate;-   forming upper and lower electrical contacts formed on the plurality    of nanowires of the predetermined region of the substrate; wherein-   the plurality of nanowires coherently emit optical radiation through    the Anderson localization of light.

In accordance with an embodiment of the invention there is provided amethod of suppressing non-radiative surface recombination within astructure formed from a compound semiconductor comprising at least apredetermined first Group III element, a predetermined second Group IIIelement, and a predetermined Group V element by establishing within apredetermined portion of the structure a region richer in concentrationof the predetermined first Group III element within a core region of thestructure than an outer region of the structure and richer inconcentration of the predetermined second Group III element within theouter region of the structure than the core region of the structure.

In accordance with an embodiment of the invention there is provided adevice comprising:

-   a substrate;-   a plurality of nanowires formed from a single compound semiconductor    over a predetermined portion of their length, the single compound    semiconductor within the predetermined portion of the nanowire    comprising at least a predetermined first Group III element, a    predetermined second Group III element, and a predetermined Group V    element, wherein within the predetermined portion of the nanowire    the predetermined first and second Group III elements have radial    profiles, wherein    -   the radial profile of the predetermined first Group III element        is richer in concentration within a core region of the nanowire        than an outer region of the nanowire; and    -   the radial profile of the predetermined second Group III element        is richer element within the outer region of the nanowire than        the core region of the nanowire; and-   upper and lower electrical contacts electrically connected to    opposing ends of the plurality of nanowires.

In accordance with an embodiment of the invention there is provided amethod comprising:

-   providing a substrate;-   growing a plurality of nanowires formed from a single compound    semiconductor over a predetermined portion of their length, the    single compound semiconductor within the predetermined portion of    the nanowire comprising at least a predetermined first Group III    element, a predetermined second Group III element, and a    predetermined Group V element, wherein within the predetermined    portion of the nanowire the predetermined first and second Group III    elements have radial profiles, wherein    -   the radial profile of the predetermined first Group III element        is richer in concentration within a core region of the nanowire        than an outer region of the nanowire; and    -   the radial profile of the predetermined second Group III element        is richer element within the outer region of the nanowire than        the core region of the nanowire; and-   forming upper and lower electrical contacts electrically connected    to opposing ends of the plurality of nanowires.

Other aspects and features of the present invention will become apparentto those ordinarily skilled in the art upon review of the followingdescription of specific embodiments of the invention in conjunction withthe accompanying figures.

BRIEF DESCRIPTION OF THE DRAWINGS

Embodiments of the present invention will now be described, by way ofexample only, with reference to the attached Figures, wherein:

FIG. 1A depicts a schematic of an AlGaN nanowire laser structureaccording to an embodiment of the invention;

FIG. 1B depicts an SEM image of self-organized AlGaN nanowire arraysformed on silicon;

FIG. 1C depicts an atomic-resolution image taken from a p-AlGaN nanowiresection grown according to an embodiment of the invention;

FIGS. 1D and 1E respectively depict line profiles across pairs ofadjacent atomic dumbbells within FIG. 1C directly confirming theN-polarity of AlGaN nanowires grown according to an embodiment of theinvention in FIG. 1C;

FIG. 1F depicts a STEM-HAADF Z-contrast image taken from a p-AlGaNregion of high Al concentration nanowire according to an embodiment ofthe invention indicating an Al-rich shell near the nanowire surface;

FIG. 1G depicts a STEM-HAADF Z-contrast image taken from an i-AlGaNregion of high Al concentration nanowire according to an embodiment ofthe invention indicating the presence of atomic-scale compositionalfluctuations.

FIG. 2A depicts room temperature photoluminescence spectra for AlGaNnanowires according to embodiments of the invention with low and high Alconcentrations;

FIG. 2B depicts the variation of internal quantum efficiency withexcitation power for AlGaN nanowires according to embodiments of theinvention with high Al concentration;

FIG. 3A depicts STEM-HAADF image of a AlGaN nanowire according to anembodiment of the invention with low Al concentration;

FIG. 3B depicts pseudo-color EELS-SI map of an AlGaN nanowire accordingto embodiments of the invention with low Al concentration showing Ga andAl elemental distributions;

FIG. 3C depicts a high-magnification image from the boxed region in FIG.3A showing uniform Ga and Al distribution in the nanowire bulk region;

FIG. 3D depicts a high-magnification image from a p-AlGaN nanowireaccording to an embodiment of the invention with high Al concentration;

FIG. 3E depicts a line profile from the boxed region in FIG. 3Ddepicting the (brighter) AlGaN quantum dash-like nanostructures;

FIG. 4A depicts schematically the formation of AlGaN nanowires accordingto embodiments of the invention;

FIG. 4B depicts schematically the fabrication procedure for AlGaNnanowire laser devices according to embodiments of the invention;

FIG. 5A depicts I-V characteristics of AlGaN nanowire lasers accordingto embodiments of the invention together with simulated electrical fielddistribution λ=290 nm;

FIG. 5B depicts electroluminescent spectra measured under differentinjection currents for AlGaN nanowire lasers according to embodiments ofthe invention;

FIG. 5C depicts electroluminescent light versus current for λ=289 nmlasing peak and spontaneous emission for AlGaN nanowire lasers accordingto embodiments of the invention;

FIGS. 5D and 5E depict the linewidth and peak wavelength as a functionof the injection current density, respectively, for AlGaN nanowirelasers according to embodiments of the invention;

FIG. 5F depicts the probability of forming cavities of Q higher than1000 for AlGaN nanowire lasers according to embodiments of theinvention;

FIG. 6A depicts an SEM image of an AllnGaN core-shell nanowire arraygrown on a silicon substrate according to an embodiment of the inventionwith structural schematic shown in the inset;

FIG. 6B depicts normalized photoluminescence spectra for AllnGaNcore-shell nanowires according to embodiments of the invention undervarying growth conditions;

FIG. 7 depicts time-resolved photoluminescent emission of threerepresentative AllnGaN core-shell nanowire LED heterostructuresaccording to embodiments of the invention, exhibiting carrier lifetimeof 0.35 ns, 0.96 ns and 1.86 ns;

FIG. 8A depicts an HAADF-STEM image for the AlInGaN core-shell LEDnanowire with shell thickness ˜18.5 nm with a peak emission of λ=495 nm;

FIGS. 8B and 8C depict elemental profiles derived from EDXS linescanning analysis of the AlInGaN core-shell nanowire in Figure A inaxial and radial directions respectively;

FIG. 9A depicts current-voltage characteristics of a representativeAlInGaN nanowire LED on Si,

FIG. 9B depicts normalized electroluminescent spectra for AlInGaNcore-shell nanowires LEDs grown under different conditions;

FIG. 9C depicts the optical output versus injection current for ˜1 mm²AlInGaN nanowire array under pulsed biasing conditions;

FIG. 9D depicts variations of the optical output of AlInGaN core-shellnanowires LEDs measured under injection currents of 90 Acm⁻² versuscarrier lifetime of the nanowire heterostructure.

DETAILED DESCRIPTION

The present invention is directed to semiconductor nanowire devices andmore particularly to ultraviolet emitters, self-organized nanoclusters,Anderson localization, and self-organized elementally rich shells.

The ensuing description provides exemplary embodiment(s) only, and isnot intended to limit the scope, applicability or configuration of thedisclosure. Rather, the ensuing description of the exemplaryembodiment(s) will provide those skilled in the art with an enablingdescription for implementing an exemplary embodiment. It beingunderstood that various changes may be made in the function andarrangement of elements without departing from the spirit and scope asset forth in the appended claims.

A. Ultraviolet Nanowire Light Source

A1. Fabrication

The inventors grew catalyst-free AlGaN nanowires directly onto Sisubstrates using a radio frequency plasma-assisted molecular beamepitaxy (MBE) system under nitrogen rich conditions. The Si wafers werecleaned by standard solvents prior to loading into the system afterwhich the Si surface oxide was thermally desorbed at 780° C. in situ.Before growth initiation, a thin Ga seeding layer was utilized topromote the nanowire formation. Schematically shown in FIG. 1A, thegrown nanowire structure consists of Si-doped GaN (˜250 nm), AlGaN p-i-njunction (each layer ˜100 nm), and a thin (˜10 nm) Mg-doped GaN contactlayer. The growth conditions include a nitrogen flow rate at 1.0standard cubic centimeter per minute (sccm), a forward plasma power of350 W, Ga beam equivalent pressure (BEP) in the range of1×10⁻⁸≤Ga_(BEP)≤6×10⁻⁸ Torr, and Al BEP in the range of1×10⁻⁸≤Al_(BEP)≤4×10⁻⁸ Torr. The growth temperatures for GaN and AlGaNlayers were 780° C. and 800° C., respectively. The Si and Mgconcentrations for n- and p-AlGaN regions were estimated to be ˜1×10¹⁸cm⁻³ and ˜1×10²⁰ cm⁻³. The typical SEM image taken with a 45-degreeangle of such AlGaN nanowires is shown in FIG. 1B. It is seen that thenanowires possess great size uniformity in terms of the nanowire lengthand diameter.

The crystal polarity of the wurtzite-structured AlGaN nanowires observedalong the <1120> orientation (a-plane) was directly determined from theHAADF image intensity across adjacent atomic dumbbells within the highAl-content AlGaN, such as depicted in FIG. 1C. The low-Z N-atoms arevisible in the proximity of the weaker scattering Al-atoms in comparisonto the strong scattering from Ga-atoms in GaN. A detailed investigationof the atomic structure in FIG. 1C shows elongation of the group IIIcolumn alternating diagonally left and right. The corresponding lineprofiles of the ADF intensity from the pair of marked atomic dumbbellsin FIG. 1C are displayed in FIGS. 1D and 1E, confirming that thepresented AlGaN nanowires are N-polar.

Now referring to FIG. 1F there is depicted a low magnificationZ-contrast image of p-AlGaN region of a high Al concentration nanowirealong the <1100> orientation (m-plane). It is seen that besides thecompositional modulations as described in the main text, a few nm thickAl-rich AlGaN shell can be seen (darker region) near the surface, whichis further surrounded by Ga-rich outermost shell. This is similar tothat observed in the low Al concentration nanowires. The atomic-scalecompositional modulations can also be observed in the i- and n-AlGaNregions in the high Al concentration nanowire. Shown in FIG. 1G is a lowmagnification Z-contrast image of the i-AlGaN region along the a-planeof a separate nanowire where it can be seen that similar compositionalmodulations also exist.

A2. Characterisation

Optical properties of AlGaN nanowires grown according to embodiments ofthe invention were studied using temperature variable photoluminescence(PL) spectroscopy. Show in FIG. 2A are the photoluminescence (PL)spectra measured at room-temperature for two representative nanowiresamples, including a low Al concentration (10% by the nominal BEP ratio)and high Al concentration sample (70% by the nominal BEP ratio), denotedas Samples A and B, respectively. Sample A was excited by a 266 nmdiode-pumped solid state laser, while sample B was excited by a 193 nmArF excimer laser. The collected PL emission was spectrally resolved anddetected by a photomultiplier tube detector. Compared to Sample A, thespectral linewidth of Sample B is about twice as broad, which can beascribed to the presence of atomic-scale compositional modulations,i.e., quantum-dot/dash-like nanostructures in AlGaN nanowires (asdescribed and discussed below). The luminescence efficiency was furtherestimated by comparing the integrated PL intensity at room-temperaturewith that measured at 10 K, assuming the luminescence efficiency at 10 Kto be unity. In practice, it is noted that the luminescence efficiencyat 10 K also varies with the excitation power and other factors;therefore, it is necessary to examine the efficiency over a broad rangeof excitation powers. Shown in FIG. 2B is the internal quantumefficiency (IQE) of AlGaN nanowires according to embodiments of theinvention as a function of excitation power for Sample B. It can be seenthat the estimated efficiency stays nearly constant in the range of70-80% over two orders of magnitude variations in the excitation power.This luminescence efficiency is significantly higher than previouslyreported AlGaN nanowires, which can be largely ascribed to the strongcarrier confinement due to the presence of quantum dot/dash-likenanostructures (as described and discussed below).

Detailed structural characterization of Samples A and B were performedby high-resolution scanning transmission electron microscope (STEM).FIG. 3A shows the high-angle annular dark-field (HAADF) Z-contrast imageof a single AlGaN nanowire from Sample A taken with a double aberrationcorrected STEM under 200 kV. It is seen that the nanowire has a lengthof ˜600 nm and a top diameter of 50 nm, with an inversely taperedmorphology. Electron energy-loss spectroscopy spectrum imaging (EELS-SI)was further performed to study the elemental distribution. Lateralprofile of Ga-map suggests that a hexagonal cross-section with m-planefacets maintained all along the nanowire. Shown in FIG. 3B, apseudo-color image of Ga- and Al-signals suggests a low Al concentrationin the core region and a high Al concentration in the shell region. Theformation of such core-shell nanowire structures can significantlysuppress non-radiative surface recombination and increase the carrierinjection efficiency of nanowire LEDs and lasers. Shown in FIG. 3C is ahigh-resolution image taken from the marked region in FIG. 3A. Theaforementioned Al-rich AlGaN shell near the nanowire surface (darkregion) can be observed; in addition, this Al-rich AlGaN shell isfurther surrounded by a few atomic-plane thick Ga-rich AlGaN outermostshell. One other important feature to point out is, for such a low Alconcentration, the nanowire bulk region shows highly uniformcompositional distribution.

With increasing Al concentration, the presence of Al-rich AlGaN shellnear the surface region can also be observed (as described and depictedin respect of FIGS. 1E and 1F); however, in contrast to Sample A (low Alconcentration sample), strong atomic-scale compositional fluctuationswere formed within the AlGaN layers in Sample B (high Al concentrationsample). Illustrated in FIG. 3D is a high-magnification Z-contrast imagefrom the p-AlGaN region of one AlGaN nanowire from Sample B. Thepresence of extensive atomic-scale Ga-/Al-rich modulations along thegrowth direction can be clearly observed. In addition, the Ga-rich AlGaN(brighter region) is not continuous along the lateral direction(perpendicular to the growth direction). Such Ga-rich AlGaN regions havesizes varying from a single atomic layer (˜0.25 nm) to 2 nm along thegrowth direction, and lateral sizes varying from 2 to 10 nm. A lineprofile from the boxed region in FIG. 3D is shown in FIG. 3E. Thelocalized variations in the Ga-concentration are estimated to be 5-10 at%. Moreover, both m- and a-plane orientation views of the Ga-rich AlGaNregions suggest that the Ga-rich regions make up a large portion of theprojected thickness within the nanowire diameter (in order to bedetectable), which, together with the observed lateral discontinuity,indicates that the atomic-scale fluctuations possess quantumdot/dash-like characteristics. It is further noted that suchatomic-scale compositional modulations can be also observed in i- andn-AlGaN regions (as described and depicted in respect of FIG. 1F).

The formation process of such quantum dot/dash-like features is relatedto the interplay between spontaneous chemical ordering and anisotropicatom migration from the irregular top/lateral surfaces of nanowirearrays. Previously, the spontaneous formation of Al-rich and Ga-richlayers was observed and explained by the significantly different bindingenergies between Ga—N and Al—N. Such spontaneous chemical orderingalone, however, cannot explain the formation of quantum dot/dash-likenanostructures. The inventors show through the embodiments of theinvention that, due to the random nucleation and formation process aswell as the shadowing/coalescence effect of neighboring nanowires,self-organized AlGaN nanowires tend to develop into non-symmetricshapes, which strongly affects the diffusion of Al and Ga atoms alongthe sidewalls during the subsequent growth process. This effect,together with the difference in surface migration rates between Ga andAl atoms, can strongly modulate the spontaneous chemical orderingprocess at the growth front. As such, quantum dot/dash-likenanostructures are formed in self-organized AlGaN nanowire arrays,illustrated in FIG. 4A. Effects of such compositional variations on theoptical properties of AlGaN nanowires are manifested by the broad PLspectral linewidth as well as the extremely high luminescence efficiency(˜80%) at room temperature, shown in FIGS. 2A and 2B. Due to the largeeffective mass of charge carriers in Al-rich AlGaN, the Bohr radii areonly 1-2 nm, which is comparable to the size variations of the observedquantum dots and dashes. As a consequence, they can provide strong 3Dquantum confinement. Moreover, such local compositional variations alsoinduce strong perturbation to the energy band, due to the changes in thepolarization fields.

Within the prior art it has been shown that such random nanowire arrayscan function as a high Q optical cavity, due to the Andersonlocalization of light. Optically pumped lasing has been realized in GaNnanowires utilizing such random cavities. More recently, the inventorshave demonstrated electrically-injected AlGaN nanowire lasers operatingin the UV-AII band. The inventor's analysis suggests that such AlGaNnanowires, with optimized size and density, can lead to strong opticalconfinement in the deep UV bands (as described and depicted in respectof FIG. 5F). The optical cavity is formed due to the multiple scatteringprocesses in the randomly distributed AlGaN nanowire arrays. Thevertical optical confinement is further made possible by the inverselytapered geometry of nanowires, as illustrated in FIGS. 1 and 3A. Shownin the inset of FIG. 5A is the simulated electric field distribution ofconfined photons (λ=290 nm) in the lateral dimension of near-randomlydistributed AlGaN nanowires (as described and depicted in respect ofFIG. 5F and Section A3).

A3. Light Confinement in Self-Organized AlGaN Nanowire Arrays

Previously, lasing phenomena has been observed in disordered nanowirearrays of various material systems. Strong light localization becomespossible as photons are recurrently scattered among disorderednanowires, due to the interference of scattered waves. The geometryparameters of nanowires play important roles in achieving effectiveoptical confinement.

The presence of any possible modes around λ=290 nm was calculated by theRF module of Comsol Mulitphysics software. The simulation was performedby varying the average nanowire diameters (d) and filling factors (F)repetitively over a wide range. In each simulation, the nanowirepositions were generated randomly, and their sizes were also randomlyvaried in the range of 0.85 d≤ϕ≤1.15 d. Such considerations were basedon the SEM measurements of self-organized AlGaN nanowire arrays on Si.The probability of finding high quality-factor (Q) modes at λ=290 nm wasthen investigated by the Monte Carlo method. Referring to FIG. 5F,effective optical confinement can be readily achieved for nanowirediameters in the range of 60 nm≤d≤75 nm and filling factors in the rangeof 15%≤F≤55%. For comparison, the average diameters of self-organizedAlGaN nanowires are d˜65 nm, and the average filling factor is F˜30%taking into account the inversely tapered nanowire geometry.Self-organized AlGaN nanowires with such size distributions wereachieved by carefully optimizing the growth conditions, including thesubstrate temperature, growth rate, and nitrogen flux. Additionally,optical confinement along the vertical direction was made possible bythe effective index guiding of the AlGaN nanowire arrays along thegrowth direction. Due to the inversely tapered geometry, the effectiverefractive index reaches minimum at the bottom of nanowires. The fillingfactor varies approximately from d˜0.07 to d˜0.7 from the bottom to themiddle of AlGaN nanowires.

A4. Laser Device Fabrication

Electrically injected laser devices were fabricated by the processdescribed below and as depicted in respect of FIG. 4B. Firstly, Ti[10nm]:Au[30 nm] n-metal layers were deposited onto the backside of n-Sisubstrate with an e-beam evaporator wherein the native silicon oxide onthe backside of n-Si wafer had been removed by HF before metaldeposition. Secondly, the sample was patterned into devices withdifferent sizes by optical lithography. In addition, in contrast toprior reported self-organized light emitting devices fabricated by theinventors. No filling materials were utilized to avoid any lightabsorption in the UV-B spectral range. Accordingly, air holes existamongst the nanowires. The top metal contact was deposited also bye-beam evaporation, with a tilting angle. Before metal deposition, thenanowire surface was etched by HCl solution to remove any surface oxide.The metallization for the top metal contact being Ni[10 nm]:Au[5 nm].Rapid thermal annealing was carried out at 550° C. for 1 minute.Photoresist patterning was employed to define the mesa emission regionsprior to top metal deposition such that device size was ˜10 μm²

The devices were characterized under continuous-wave (CW) operation. TheI-V characteristics measured at room temperature are shown in FIG. 5Awhere it can be seen that an excellent diode was formed with negligibleleakage current. The device has a turn on voltage around 5 V. Theexcellent current-voltage characteristics measured, compared toconventional planar devices, is due to the significantly enhanced dopantincorporation in the near-surface region of nanowire structures and theresulting efficient charge carrier hopping conduction. Theelectroluminescence (EL) spectra were measured at room temperature fromthe nanowire top surface. The output light was collected by a deep UVobjective, spectrally resolved by a high-resolution spectrometer, anddetected by a liquid nitrogen cooled CCD.

The EL spectra under different injection currents are shown in FIG. 5B.It is seen that at low injection currents (˜100 A/cm2), only a veryweak, broad emission spectrum (black curve) can be measured. As theinjection current increases, an emission peak centered around 289 nmappears. It increases rapidly with injection current and becomesdominant at relatively high current densities. FIG. 5C shows theintegrated EL intensity of the 289 nm peak vs. the injection current,which exhibits a clear threshold at ˜300 A/cm2. Variations of thebackground emission with increasing current are also shown in FIG. 5C(black triangle), which was taken from the boxed region (spectral width˜3 nm) in FIG. 5B. Compared to the lasing peak at 289 nm, the integratedbackground emission shows a negligible increase above threshold, whichis explained by the clamping of carrier concentration above threshold.The inset of FIG. 5C shows the L-I curve of the lasing peak at 289 nm ina logarithmic scale. Three distinct regions including spontaneousemission, amplified spontaneous emission, and lasing emission can beclearly observed, further providing an unambiguous evidence for theachievement of lasing. The lasing threshold of 300 A/cm2 issignificantly smaller compared to the previously reported electricallyinjected quantum well lasers in the UV-A band (20 kA/cm2 at 336 nm). Theextremely low lasing threshold is attributed to the drastically reducedtransparency carrier density of 3D quantum-confined nanostructures, thenearly defect-free AlGaN core-shell nanowires, and the high Q opticalcavity offered by Anderson localization.

The derived linewidth versus injection current density is shown in FIG.5D. It can be seen that as the injection current increases, thelinewidth decreases from 6 nm below threshold to 2.6 nm above threshold,further confirming lasing action. The relatively broad lasing linewidth(2.6 nm) is likely related to the presence of multiple modes in therandom cavity and the broad gain spectrum due to the size dispersion ofquantum dots/dashes. Lasing wavelength versus injection current densitywas also investigated, illustrated in FIG. 5E. The wavelength is nearlyinvariant above threshold, suggesting stable exciton emission. Thedevice output power is estimated to be in the range of which is largelylimited by the light absorption of the p-GaN and p-metal contact layers.

Accordingly, the inventors have demonstrated spontaneous formation ofquantum dot/dash-like nanostructures in self-organized AlGaN nanowirearrays which through their resulting 3D quantum confinement, togetherwith the nearly defect-free nanowire structures, drastically reduce thecurrent density required for population inversion, leading toelectrically injected AlGaN nanowire lasers with relatively lowthreshold at room-temperature. In this manner an electrically injectedUV laser (λ˜290 nm) has been established facilitating development ofelectrically injected small-scale deep UV lasers.

B: Self-Organizing Core-Shell Nanowire Heterostructures

B1: Growth

The inventors grew catalyst-free AlInGaN core-shell nanowire LEDheterostructures on Si(111) substrates using a radio frequencyplasma-assisted MBE system under nitrogen-rich conditions as previouslyestablished by the inventors for growing high quality nanowires withoutforeign metal catalysts. Schematically as shown in the inset of FIG. 6A,the AlInGaN nanowire LED heterostructure consists of a lower segment of˜200 nm GaN:Si, 70 nm AlInGaN central region, and 150 nm GaN:Mg uppersegment. During the growth, the nitrogen flow rate was kept at 1standard cubic centimeter per minute (SCCM), with a forward power of 350W. The GaN:Si and GaN:Mg segments were grown at ˜700° C. For the AlInGaNactive region, the substrate temperature was varied in the range of 620°C.≤T≤700° C. Also during the active region growth, the Ga beamequivalent pressure (BEP) was kept at 4.5×10⁻⁸ Torr, whilst the In andAl BEP were varied from 1.2×10⁻⁷ Torr≤BEP_(In)≤2.6×10⁻⁷ Torr and2.9×10⁻⁹ Torr≤BEP_(Al)≤1.3×10⁻⁸ Torr respectively.

B2: Analysis

Under growth conditions described above in Section B1 thephotoluminescence (PL) emission wavelengths of AlInGaN quaternarynanowire structures can be continuously varied across nearly the entirevisible spectral range, from 410 nm≤λ_(PL)≤630 nm as depicted in FIG.6B. Moreover, through extensive studies and variations in the growthconditions by the inventors it was also found that the Al-rich shellthicknesses of AlInGaN quaternary nanowire heterostructures can becontrollably varied, as described below, allowing for the directcorrelation with the carrier dynamics and device performance. Referringto FIG. 1A there is depicted a scanning electron microscope (SEM) imageof AlInGaN core-shell nanowires, which are vertically aligned on theSi(111) substrate. The nanowires exhibit a high degree of sizeuniformity, with the top sizes in the range of ˜70-80 nm.

Within the following section the inventors describe the carrier dynamicsand structural properties of three representative nanowire LEDheterostructures, referred to as Samples A, B, and C, with peak emissionwavelengths of 495-515 nm. The corresponding growth conditions andoptical characterization results for these three representative nanowireLED heterostructures are described in Table 1. Growth temperatures of625° C., 635° C. and 670° C. were used for the active regions of SamplesA, B, C, respectively. Compared to Sample A, the A1 BEP of Sample B wasincreased from 3.54×10-9 to 5.70×10−9 Torr. Relatively higher In and A1flux was utilized for Sample C in order to achieve enhanced shellcoverage.

TABLE 1 Growth Conditions for Representative AlInGaN Segments BEP_(Al)BEP_(Im) × 10⁻⁷ λ Symbol (×10⁻⁹ Torr) Torr T_(Sample) (° C.) (nm)τ_(Carrier) (ns) A 3.54 1.25 625 515 0.35 B 5.70 1.38 635 515 0.96 C10.7 2.06 670 495 1.86

Time-resolved PL measurements were performed to study the carrierdynamics of AlInGaN core-shell nanowire heterostructures. A pulsed 375nm diode laser with a 100 MHz repetition rate was employed as theexcitation source, which was focused on the sample surface through a 50×objective lens. The signal was detected by a photon counter with a λ>400nm long pass filter. The carrier life time τ_(Carrier) was then derivedby a standard stretched exponential model given by Equation (1) whereI(t) is the PL intensity as a function of time, and n is the stretchingparameter. Shown in FIG. 7 are the representative curves measured forsamples A, B, and C, with the derived carrier lifetimes ρ_(Carrier)=0.35ns, Σ_(Carrier)=0.96 ns, and τ_(Carrier)=1.86 ns, respectively.

$\begin{matrix}{{I(t)} = {{I(0)}e^{{- {(\frac{t}{\tau})}}n}}} & (1)\end{matrix}$

Accordingly, the inventors established that measured carrier lifetimesof AlInGaN LED heterostructures vary dramatically depending on thenanowire growth conditions. For comparison, the carrier lifetime of atypical InGaN/GaN nanowire LED structure without the incorporation ofany A1 was measured to be in the range of 0.2 ns, dominated bynon-radiative surface recombination. The carrier lifetime of suchInGaN/GaN nanowire LED structure was enhanced by the inventors to ˜0.4ns by adding an AlGaN shell. It is worth mentioning that the carrierlifetime remained nearly invariant when the excitation power was changedby over two orders of magnitude in the present study.

In order to further identify the correlation between the carrierdynamics of AlInGaN nanowire LED heterostructures and the associatedgrowth conditions, the inventors performed extensive structuralcharacterization by scanning transmission electron microscope (STEM).The high-angle annular dark-field (HAADF) atomic-number contrast imageof a nanowire from Sample C (carrier lifetime τ_(Carrier)=1.86 ns) isshown in FIG. 8A. The HAADF image shows bright contrast at the centerand dark contrast at the sidewall of the AlInGaN segment, indicating theformation of a core-shell structure. The energy dispersive x-rayspectrometry (EDXS) analysis are shown in FIG. 8B, for Ga, Al, In, andN; and FIG. 8C for Ga, Al, and In. The elemental profiles along thegrowth direction (line a-b in FIG. 8A) as depicted in FIG. 8B presentclearly one AlInGaN segment in the middle region and two GaN segments inthe top and bottom regions of the nanowire. The tail of Al profile(position: 90-130 nm) originates from the direct impingement anddiffusion of Al atoms at the sidewall of the bottom GaN:Si segmentduring the growth of AlInGaN active region. In FIG. 8C the EDXS axialline-scan analysis along the lateral dimension of the nanowire activeregion (line c-d in FIG. 8A) is presented. Here it can be seen that theAl and Ga signals extend over a much wider range than the In signal,providing unambiguous evidence that an AlGaN shell is formed surroundingan In-rich core region. The AlGaN shell thickness was estimated to be˜18.5 nm.

Similar STEM studies including HAADF and EDXS analysis were conducted onother nanowire samples. As a result, the inventors identified a directcorrelation between the formation of an Al-rich shell structure and thecarrier lifetime. For example, for Sample A, a small Al-rich shell wasformed, due to the use of very low Al flux. Accordingly, the carrierlifetime for this sample was measured to be τ_(Carrier)=0.35 ns, limitedby non-radiative surface recombination. The AlGaN shell thicknesses wereestimated to be ˜13 nm for Sample B, which leads to enhanced carrierlifetime τ_(Carrier)=0.96 ns, due to the reduced non-radiative surfacerecombination. The inventor's extensive structural and opticalcharacterizations of AlInGaN nanowire heterostructures, not presentedhere, confirmed that the carrier lifetime increased with increasingAlGaN shell thickness.

The spontaneous formation of core-shell nanowire structures can beexplained by the differences in the diffusion and desorption processesof In, Ga, and Al adatoms during the growth of the AlInGaN segment. Atelevated growth temperatures, In atoms experience much strongerdesorption compared to Ga and Al atoms on the nanowire lateral surfaces.The desorbed In atoms cannot be immediately compensated by impingingatoms, due to the shadowing effect of neighboring nanowires. As aconsequence, an Al-rich shell is formed surrounding an In-rich coreregion. It is also evident that properties of the core-shell structurescan be controlled by varying the substrate temperature, in addition tothe group III/N flux ratios, since the sticking and diffusioncoefficients of atoms are very sensitive to the substrate temperature.

In order to examine the impact of Al-rich shell structure on the deviceperformance, we have fabricated and characterized large area AlInGaNnanowire LEDs with areal sizes of ˜0.3×0.3 mm² to ˜1.0×1.0 mm². Thedevice fabrication process involves the use of polyimide surfacepassivation and planarization, standard photolithography and contactmetallization techniques. As depicted by the insert in FIG. 6A, p- andn-contacts were deposited on the GaN:Mg and the Si substrate backside,respectively.

Output characteristics of ˜1.0×1.0 mm² AlInGaN core-shell LEDs weremeasured under various injection currents under pulsed biasingconditions (1% duty cycle) to minimize junction heating effect. Shown inFIG. 9A is a representative I-V curve of AlInGaN core-shell nanowireLEDs fabricated by the inventors according to embodiments of theinvention, which exhibit excellent current-voltage characteristics. Themeasured electroluminescence (EL) spectra of various AlInGaN nanowireLEDs are shown in FIG. 9B. Tunable emissions from 410nm≤λ_(Emission)≤630 nm can be readily achieved. Variations of the outputpower versus injection current density for a few representative LEDstructures are further shown in FIG. 9C. It is evident thatsignificantly higher output power can be achieved for nanowire LEDs withlonger carrier lifetime, due to the suppression of non-radiative surfacerecombination, e.g. comparing Sample C to Sample A. Under an injectioncurrent density of 100 Acm⁻², an output power of >30 mW was measured forSample C, which is significantly higher than that of previously reportedaxial InGaN nanowire LEDs within the prior art.

Variations of the measured output power at an injection current densityof 90 Acm⁻² versus carrier lifetime are further summarized in FIG. 9Dfor LEDs operating in the blue-green wavelength range (410nm≤λ_(Emission)≤530 nm). It is seen that the output power increasesdramatically with carrier lifetime, which is also consistent with recentreports on AlGaN/InGaN dot-in-a-wire LEDs by the inventors. Such studiesprovide unambiguous evidence that robust core-shell structures areextremely beneficial in suppressing any non-radiative surfacerecombination and to achieve high power operation for axial nanowiredevices. The direct correlation between carrier lifetime and deviceperformance also rules out the possibility that the enhanced carrierlifetime is caused by the quantum-confined Stark effect (QCSE) withincreasing Al incorporation, since QCSE generally leads to reducedefficiency and output power.

Accordingly, the inventors have demonstrated full-color AlInGaNcore-shell quaternary nanowire LEDs grown directly onto Si substrateswherein through controlled growth parameters the spontaneous formationof an In-rich core and Al-rich shell structure occurs establishing aradial carrier confinement which can suppress the undesirablenon-radiative surface recombination, leading to enhanced carrierlifetime and significantly increased output power.

Whilst the principle has been established in respect of AlInGaNnanowires other compound semiconductor nanowire structures may beengineered to yield such a radial carrier confinement therebysuppressing undesirable non-radiative surface recombination at thenanowire surface.

The foregoing disclosure of the exemplary embodiments of the presentinvention has been presented for purposes of illustration anddescription. It is not intended to be exhaustive or to limit theinvention to the precise forms disclosed. Many variations andmodifications of the embodiments described herein will be apparent toone of ordinary skill in the art in light of the above disclosure. Thescope of the invention is to be defined only by the claims appendedhereto, and by their equivalents.

Further, in describing representative embodiments of the presentinvention, the specification may have presented the method and/orprocess of the present invention as a particular sequence of steps.However, to the extent that the method or process does not rely on theparticular order of steps set forth herein, the method or process shouldnot be limited to the particular sequence of steps described. As one ofordinary skill in the art would appreciate, other sequences of steps maybe possible. Therefore, the particular order of the steps set forth inthe specification should not be construed as limitations on the claims.In addition, the claims directed to the method and/or process of thepresent invention should not be limited to the performance of theirsteps in the order written, and one skilled in the art can readilyappreciate that the sequences may be varied and still remain within thespirit and scope of the present invention.

What is claimed is:
 1. A device comprising: a substrate; a plurality ofnanowires formed from a single compound semiconductor over apredetermined portion of their length, the single compound semiconductorwithin the predetermined portion of the nanowire comprising at least apredetermined first Group III element, a predetermined second Group IIIelement, and a predetermined Group V element, wherein within thepredetermined portion of the nanowire the predetermined first and secondGroup III elements have radial profiles, wherein the radial profile ofthe predetermined first Group III element is richer in concentrationwithin a core region of the nanowire than an outer region of thenanowire; and the radial profile of the predetermined second Group IIIelement is richer element within the outer region of the nanowire thanthe core region of the nanowire; and upper and lower electrical contactselectrically connected to opposing ends of the plurality of nanowires.2. The device according to claim 1, wherein the radial profile of thepredetermined first Group III element and the radial profile of thepredetermined second Group III element provide a core-shell geometrywithin the predetermined portion of the length of the nanowire andcontinuously vary having been formed spontaneously forms during growthof the nanowire.
 3. The device according to claim 1, wherein theplurality of nanowires are nanowire light emitting diodes wherein withinthe predetermined portion non-radiative surface recombination is reducedwith respect to that of the single compound semiconductor without theradial profiles resulting in an increased carrier lifetime.
 4. Thedevice according to claim 1, wherein the carrier lifetime within thepredetermined portion of the nanowire is at least one nanosecond.
 5. Thedevice according to claim 1, wherein the predetermined first Group IIIelement is indium; the predetermined second Group III element isaluminum; the predetermined Group V element is nitrogen; and thenanowires further comprise gallium and under forward electrical bias arelight emitting diodes wherein their emission wavelength is determined bythe relative proportions of the aluminum, indium, and gallium.
 6. Amethod comprising: providing a substrate; growing a plurality ofnanowires formed from a single compound semiconductor over apredetermined portion of their length, the single compound semiconductorwithin the predetermined portion of the nanowire comprising at least apredetermined first Group III element, a predetermined second Group IIIelement, and a predetermined Group V element, wherein within thepredetermined portion of the nanowire the predetermined first and secondGroup III elements have radial profiles, wherein the radial profile ofthe predetermined first Group III element is richer in concentrationwithin a core region of the nanowire than an outer region of thenanowire; and the radial profile of the predetermined second Group IIIelement is richer element within the outer region of the nanowire thanthe core region of the nanowire; and forming upper and lower electricalcontacts electrically connected to opposing ends of the plurality ofnanowires.
 7. The method according to claim 6, wherein the radialprofile of the predetermined first Group III element and the radialprofile of the predetermined second Group III element provide acore-shell geometry within the predetermined portion of the length ofthe nanowire which spontaneously forms during growth of the nanowire. 8.The method according to claim 6, wherein the plurality of nanowires arenanowire light emitting diodes wherein within the predetermined portionnon-radiative surface recombination is reduced with respect to that ofthe single compound semiconductor without the radial profiles resultingin an increased carrier lifetime.
 9. The method according to claim 6,wherein the carrier lifetime within the predetermined portion of thenanowire is at least one nanosecond.
 10. The method according to claim6, wherein the predetermined first Group III element is indium; thepredetermined second Group III element is aluminum; the predeterminedGroup V element is nitrogen; and the nanowires further comprise galliumand under forward electrical bias are light emitting diodes whereintheir emission wavelength is determined by the relative proportions ofthe aluminum, indium, and gallium.
 11. The method according to claim 6,wherein the radial profiles of the predetermined first and second GroupIII elements are established through self-organization during growth ofthe plurality of nanowires.